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Oxidation of GeSi
MICROELECTRONIC ENGINEERING ELSEVIER
Microelectronic
Engineering
28 (1995) 185-191
Oxidation of GeSi Marc-A. Nicoleta and Wen-Shu Liub aCalifornia Institute of Technology, Pasadena, CA 91125, USA blntegrated Device Technology, Inc., San Jose, CA 95134, USA
Deal and Grove formulated a model for the oxidation of silicon in 1965 that has served since as a basis for all subsequent work on the subject [l]. The model assumes that it is the oxidant that moves through the oxide and forms new oxide by reaction with the silicon at the silicon/oxide interface. This model leads to the equation X*
+
AX = B(t - ti),
(1)
where x is the oxide thickness, t is the oxidation duration, and A and 6 are constants. The quantity ti = (Xi* + AXi)/B is related to the initial condition, where Xi is the thickness of a preexisting oxide. Two limiting growth regimes exist: (i) when t c-c A*/4B, the reaction mechanism at the interface limits the growth rate and it becomes linear, x = (B/A)(t - ti), and (ii) when t >> A*/4B, the diffusion of the oxidant through the oxide controls the growth rate and it becomes parabolic, x2 = Bt. The activation energy of the diffusion rate constant B (about 0.7 and 1.2 eV for wet and dry ambients, respectively) is similar to that for diffusion of oxidants through fused quartz, while the interfacial rate constant B/A is the same for wet and dry conditions and has an activation energy close to the energy needed to break a Si-Si bond (approximately 2 eV). The oxidation of GeSi is more complex than that of Si because the binary Si-0 system turns into a ternary one when germanium is added. Beginning with the very early publications [2], most investigations of GeSi oxidation have been confined to Ge-poor alloys (i.e. to x c 0.2 for GexSil_x) and to epitaxial layers grown on singlecrystalline substrates of silicon. These are the conditions of interest for Si-based heteroepitaxial devices. Unless otherwise stated, what follows applies to such epitaxial films of GeSi. The complexity that comes with the addition of germanium is clearly demonstrated by the fact that amorphous oxides of different compositions can be obtained with GeSi, depending on the oxidation conditions. By contrast, with pure silicon, changes in temperature, ambient (wet or dry), or pressure alter the growth kinetics, but amorphous silicon dioxide is always what forms. The main parameter controls the germanium and silicon content in the oxide of GeSi is the temperature, 0167.9317/95/$09.50 D 1995 - Elsevier Science B.V. All rights reserved. SSDI 0 167-93 17(95)00040-2
186
M.A. Nicolet,
W.S. Liu / Microelectronic
Engineering
28 (1995) 185-191
but the alloy composition, and the nature and pressure of the oxidizing ambient also affect the outcome. For example, wet oxidation at 900°C or above and at atmospheric pressure of GeSi with less than about 20”at. % germanium produces a layer of silicon dioxide that is essentially bare of germanium dioxide. The oxide layer, however, grows faster on GeSi than on silicon. It is now known that this effect is due to Ge which speeds up the reaction process at the alloy/dioxide interface, that is, the value of the parameter A is reduced. The effect has been explained variously as a consequence of the weaker Ge-Si bond compared to the Si-Si bond, a catalytic action of Ge at the interface, or the ability of germanium to suppress the generation of interstitials in the alloy. Since only silicon dioxide forms at elevated temperatures, the germanium content of the alloy beneath the oxide increases by in diffusion of the germanium. The epitaxial structure of the alloy is thereby preserved. Quite a different situation prevails at 600 or 700°C and wet oxidation. The oxide that forms on a Ge,Sit_, alloy has the same content of germanium and silicon and may be thought of as a mixture of composition (Ge02)x(Si02)t_x. The dioxide grows parabolically with the duration of oxidation almost immediately, while on Si the dioxide grows first linearly with duration for a long period. This difference again reflects a much enhanced rate of reaction (reduced value of A) for GeSi over that for silicon. When the oxide contains germanium as well, the numerical value of the parabolic rate constant B depends on the composition, B(x), but its activation energy (1 .I f 0.2 eV) does not change, at least up to x < 0.5, and is equal to that for pure silicon below 900°C (1.17 eV [3]). This means that the mechanism of transport of the oxidant in the oxide changes little with the incorporation of germanium dioxide, but that the amount of oxidant that diffuses through the oxide must increase with the germanium content while being only a weak function of temperature. A plausible interpretation of this fact is that the germanium dioxide increases the solid solubility of moisture in the mixed dioxide. This solubility is indeed much lower in SiO2 than in Ge02, which is hygroscopic. It is known that when the pressure of steam is increased, the solid solubility of moisture in Si02 rises and that therefore the flux of the oxidant through the oxide rises too. Increasing the content of germanium dioxide thus has the same effect as increasing the pressure. The activation energy remains 1.l eV in both cases. Why only silicon dioxide forms when GeSi is oxidized at elevated temperatures and germanium is expelled from the oxide into the alloy has both thermodynamic and kinetic reasons. The ternary phase diagram (Fig. l., [4]) for oxygen, germanium and silicon shows that as long as elemental silicon or GeSi alloy is still present in the system, only silicon dioxide and elemental germanium can coexist. Any germanium dioxide present in the system will be reduced to elemental germanium by the formation of additional silicon dioxide. This process can occur only as long as silicon and germanium are mobile enough in the alloy to move with respect to each other. That is evidently the case at elevated temperatures and explains why only silicon dioxide forms. The diffusivities of germanium and silicon in GeSi have an activation energy of 3 to 5 eV [5] which is several eVs above the activation energies of the oxidation process. With failing temperature,
M.A. Nicolet, W.S. Liu /Microelectronic
Engineering 28 (1995) 185-191
187
Si Ge,Si,
_x
Figure 1. Si-Ge-0 phase diagram valid up to 1000°C (from [4]). the diffusivities of germanium and silicon in the alloy thus decrease much more rapidly than those of the oxidant in the oxide or the reaction at the interface. Ai 6OO”C, the alloy is therefore kinetically frozen, but the oxidation process is not, SC, that both silicon and germanium are oxidized as the oxidant reaches the alloy interface. This is what one observes for wet oxidation. Since germanium also enhances the reaction rate at the interface, the growth quickly becomes transportlimited and quadratic in the duration of oxidation. The structure of the mixed GeSi dioxide is amorphous and contains bonds characteristic of SiO2 in tetrahedral coordination, and of Ge02 in both tetrahedral (four-fold) and octahedral (six-fold) coordination [6]. The dry oxidation of GeSi in the low temperature regime has not been studied yet because the reaction is quite slow, but it is clear that at sufficiently low temperatures germanium will be oxidized as well. The selective oxidation of elements in an alloy occurs whenever a strong non symmetry exists in their free energy change upon oxidation, when the nonoxidizing elements can diffuse easily into the bulk, and when their segregation coefficient is large. Examples are fairly common, such as transition metal silicides, CuNi alloys, or amorphous metallic ZrAINi. At low temperatures, when the growth is parabolic, the oxide grown on polycrystalline and pseudomorphic GeSi have the same thickness (same rate constant B). The microstructure of, and the strain in, the GeSi alloy thus leave the oxidation process unchanged. This fact is consistent with the finding that the oxidation process is controlled by the transport across the oxide. Reducing the content of germanium in the alloy shifts the onset of germanium rejection to lower temperature. Going from an epitaxial to a polycrystalline film does the same thing because grain boundaries enhance diffusivities. The instability of GeSi02 in the presence of pure Si or of GeSi alloy that the phase diagram of Fig. 1. predicts is experimentally verifiable even at room temperature. After several months, crystalline precipitates appear at the oxide/alloy
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M.A. Nicolet, W.S. Liu / Microelectronic Engineering 28 (1995) 185191
Figure 2. Cross-sectional transmission electron that micrographs show that a GeSi02 oxide is unstable in non with contact oxidized GeSi alloy (from [a]).
M.A. Nicolet, W. S. Liu / Microelectronic Engineering 28 (199.5) 185-191
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interface (Fig. 2.). The precipitates have a lattice constant that is similar to that of the underlying GeSi alloy. This reaction is quite rapid at elevated temperatures (Fig. 2.). An instability of this kind is unacceptable from a device point of view and means that the gate of GeSi MOSFETs may not be formed at low temperatures simply by wet oxidation of the alloy. Dry oxidation at high temperatures does produce a stable alloyBi02 interface, but the density of interfacial states is too high for practical MOSFETs. Seen from this vantage point, the oxidation of GeSi alloy holds no promise for MOS device applications in the way the direct oxidation of pure Si does. The unstable nature of GeSi02 offers interesting opportunities for controlled phase transformation, however. Fig. 3 shows a schematic of a process sequence where an epitaxial layer of a GeSi alloy on (lOO)Si is first oxidized in a wet ambient at 700°C. This first step of oxidation is terminated before the alloy is fully consumed by the oxidation reaction. The second step is a reduction in forming gas (95% N2 + 5% H2). What happens during this reduction step depends on the concentration of Ge. If the alloy and the oxide are Ge-rich, the oxide transforms into a trilayered configuration of Si02 on top, epitaxial Ge below, and a mixed Ge and Si02 layer at the bottom, adjacent to the alloy substrate (see right-most sketch of Fig. 3.). This latter layer (the “percolation” layer) is notable in two respects: (i) it is formed largely on account of the alloy substrate during the reduction reaction, and (ii) it consists of an interconnected network of single-crystalline Ge walls and pillars that reach (“percolate”) around Si02 inclusions and establish continuous paths between the underlying alloy and the overlying Ge layer. This percolation layer is itself subdivided into distinct sublayers, as the cross-sectional micrograph of Fig. 4. shows. The figure also displays one of two remarkable properties of the Ge layer: it has fewer dislocations than the GeSi alloy seed on which it grows (about 109 versus 10 ‘0 dislocations/cm2). In addition, the Ge layer also is elastically fully relaxed, i.e. fully decoupled from its alloy substrate [7].
Figure 3. Schematic of the oxidation-reduction process sequence. The changes in the layers are approximately to scale (from [7]).
190
M.A. Nicolet, W.S. Liu /Microelectronic
Engineering 28 (1995) 185-191
Figure 4. Cross-sectional transmission electron micrograph of an epitaxial Geo_+i0.18 alloy after having been subjected to the process sequence of Fig. 3 (from [7]). The top SiO2 layer has been chemically removed prior to the crosssectioning of the sample. When the alloy and the GeSi02 are not rich in Ge, a percolation layer does not form, presumably because not enough germanium is available from the reduction process to form a connected network of Ge in SiO2. Rather, the reduction process results in the nucleation of nanocrystalline Ge precipitates distributed uniformly throughout the oxide layer that now consists of SiO2. The nanocrystillites are photoluminescent. The percolation layer does not form either, even with a high concentrations of Ge in the alloy and in the oxide, if pure hydrogen is used in the reduction step instead of forming gas. The strong driving force provided by pure hydrogen ambient results again in a homogeneous nucleation of Ge precipitates. It is the heterogeneous nucleation of Ge at the oxide/alloy interface from a germanium-rich alloy and oxide that favor the development of a percolation layer and thereby the growth of a relaxed epitaxial Ge layer whose crystalline quality is superior to that of the alloy from which it originates. As the structure of the percolation layer reveals, the details of the reduction mechanisms must be complex. This structure clearly results from simultaneous and competing reactions that occur in the early stages of the reduction process. The reduction of GeO2 to Ge by hydrogen near the oxide/alloy interface creates vapor which can oxidize additional alloy that itself can be reduced again by hydrogen, and so on. In effect, hydrogen acts as a catalyst that transfers oxygen from Ge02 to Si
M.A. Nicolet, W.S. Liu / Microelectronic Engineering 28 (199.5) 185191
191
deeper into the alloy and thereby creates SiO2 and Ge side by side. The speed at which these competing reactions progress must affect the morphology of the percolation layer. In parallel to this interfacial process, there must be processes that take place in the bulk of the GeSi dioxide that account for the massive relative displacements of Ge and SiO2. A dominant reaction there is most likely the partial reduction of Ge02 to GeO, which is volatile. Evidence for its presence is also found in the loss of Ge to the ambient that is observed at the beginning of the reduction process. As germanium is lost to the ambient, a dense layer of pure SiO2 begins to form near the surface. This change at the surface can explain why the loss of germanium ceases after a while. The transport of germanium as GeO in the GeSi02 also promotes the precipitation of germanium at the inter-face with the percolation layer. The growth is seeded by the germanium in the percolation layer. As the epitaxial germanium layer expands into a continuous film, the transport of gases accross the percolation layer slows down, and so does the growth of that layer. Evidently, the percolation layer is able to block the propagation of misfit dislocations across it by virtue of the discontinuous second phase of SiO2 islands that is distributed in the germanium matrix, since most of the misfit dislocations in the alloy terminate there. The scenario depicted here is largely conjectural. It will take additional investigations to substantiate or refute it. The potential payoff seems worth it, though, because the ability to improve the crystalline quality of an epitaxial layer over that of its seed is a remarkable phenomenon. Also, the growth of a highquality unstrained single-crystalline film of germanium on silicon is a capability that would permit the combination of integrated logic circuits in silicon and integrated optical systems in GaAs on the same silicon wafer. Research on GeSi is supported by the Semiconductor Research Corporation under a coordinated program between Caltech (95SJ-100) and UCLA (95SJ088), a support that is here gratefully acknowldged. We also express our gratitude to Integrated Device Technology, Inc.,San Jose, CA, for its generous grants.
References 1. B.E. Deal, A.S. Grove, J. Appl. Phys. 36 (1965) 770. 2. F.K. LeGoues, R. Rosenberg, T. Nguyen, F. Himpsel, B.S. Meyerson, J. Appl. Phys. 65 (1989) 1724. 3. R.R. Razouk, L.N. Lie, B.E. Deal, J. Electrochem. Sot. 128 (1981)2214. 4. D.C. Paine, C. Caragianis, A.F. Schwartzman, J. Appl. Phys. 70 (1991) 5076. 5. W.S. Liu, G. Bai, M.-A. Nicolet, C.H. Chern, V. Arbet, K.L. Wang, Mat. Res. Sot. Symp. Proc. Vol. 220 (1991) 259. 6. W.S. Liu, E.W. Lee, M-A. Nicolet, V. Arbet-Engels, K.L. Wang, N.M. Abuhadba, C.R. Aita, J. Appl. Phys. 71 (1992) 4015. 7. W.S. Liu, J.S. Chen, D.Y. Lie, M.-A. Nicolet, Appl. Phys. Lett. 63 (1993) 1405. 8. W.S. Liu, J.S. Chen, M.-A. Nicolet, V. Arbet-Engels, K.L. Wang, J. Appl. Phys. 72 (1992) 4444.
Microelectronic
Engineering
28 (1995) 185-191
Oxidation of GeSi Marc-A. Nicoleta and Wen-Shu Liub aCalifornia Institute of Technology, Pasadena, CA 91125, USA blntegrated Device Technology, Inc., San Jose, CA 95134, USA
Deal and Grove formulated a model for the oxidation of silicon in 1965 that has served since as a basis for all subsequent work on the subject [l]. The model assumes that it is the oxidant that moves through the oxide and forms new oxide by reaction with the silicon at the silicon/oxide interface. This model leads to the equation X*
+
AX = B(t - ti),
(1)
where x is the oxide thickness, t is the oxidation duration, and A and 6 are constants. The quantity ti = (Xi* + AXi)/B is related to the initial condition, where Xi is the thickness of a preexisting oxide. Two limiting growth regimes exist: (i) when t c-c A*/4B, the reaction mechanism at the interface limits the growth rate and it becomes linear, x = (B/A)(t - ti), and (ii) when t >> A*/4B, the diffusion of the oxidant through the oxide controls the growth rate and it becomes parabolic, x2 = Bt. The activation energy of the diffusion rate constant B (about 0.7 and 1.2 eV for wet and dry ambients, respectively) is similar to that for diffusion of oxidants through fused quartz, while the interfacial rate constant B/A is the same for wet and dry conditions and has an activation energy close to the energy needed to break a Si-Si bond (approximately 2 eV). The oxidation of GeSi is more complex than that of Si because the binary Si-0 system turns into a ternary one when germanium is added. Beginning with the very early publications [2], most investigations of GeSi oxidation have been confined to Ge-poor alloys (i.e. to x c 0.2 for GexSil_x) and to epitaxial layers grown on singlecrystalline substrates of silicon. These are the conditions of interest for Si-based heteroepitaxial devices. Unless otherwise stated, what follows applies to such epitaxial films of GeSi. The complexity that comes with the addition of germanium is clearly demonstrated by the fact that amorphous oxides of different compositions can be obtained with GeSi, depending on the oxidation conditions. By contrast, with pure silicon, changes in temperature, ambient (wet or dry), or pressure alter the growth kinetics, but amorphous silicon dioxide is always what forms. The main parameter controls the germanium and silicon content in the oxide of GeSi is the temperature, 0167.9317/95/$09.50 D 1995 - Elsevier Science B.V. All rights reserved. SSDI 0 167-93 17(95)00040-2
186
M.A. Nicolet,
W.S. Liu / Microelectronic
Engineering
28 (1995) 185-191
but the alloy composition, and the nature and pressure of the oxidizing ambient also affect the outcome. For example, wet oxidation at 900°C or above and at atmospheric pressure of GeSi with less than about 20”at. % germanium produces a layer of silicon dioxide that is essentially bare of germanium dioxide. The oxide layer, however, grows faster on GeSi than on silicon. It is now known that this effect is due to Ge which speeds up the reaction process at the alloy/dioxide interface, that is, the value of the parameter A is reduced. The effect has been explained variously as a consequence of the weaker Ge-Si bond compared to the Si-Si bond, a catalytic action of Ge at the interface, or the ability of germanium to suppress the generation of interstitials in the alloy. Since only silicon dioxide forms at elevated temperatures, the germanium content of the alloy beneath the oxide increases by in diffusion of the germanium. The epitaxial structure of the alloy is thereby preserved. Quite a different situation prevails at 600 or 700°C and wet oxidation. The oxide that forms on a Ge,Sit_, alloy has the same content of germanium and silicon and may be thought of as a mixture of composition (Ge02)x(Si02)t_x. The dioxide grows parabolically with the duration of oxidation almost immediately, while on Si the dioxide grows first linearly with duration for a long period. This difference again reflects a much enhanced rate of reaction (reduced value of A) for GeSi over that for silicon. When the oxide contains germanium as well, the numerical value of the parabolic rate constant B depends on the composition, B(x), but its activation energy (1 .I f 0.2 eV) does not change, at least up to x < 0.5, and is equal to that for pure silicon below 900°C (1.17 eV [3]). This means that the mechanism of transport of the oxidant in the oxide changes little with the incorporation of germanium dioxide, but that the amount of oxidant that diffuses through the oxide must increase with the germanium content while being only a weak function of temperature. A plausible interpretation of this fact is that the germanium dioxide increases the solid solubility of moisture in the mixed dioxide. This solubility is indeed much lower in SiO2 than in Ge02, which is hygroscopic. It is known that when the pressure of steam is increased, the solid solubility of moisture in Si02 rises and that therefore the flux of the oxidant through the oxide rises too. Increasing the content of germanium dioxide thus has the same effect as increasing the pressure. The activation energy remains 1.l eV in both cases. Why only silicon dioxide forms when GeSi is oxidized at elevated temperatures and germanium is expelled from the oxide into the alloy has both thermodynamic and kinetic reasons. The ternary phase diagram (Fig. l., [4]) for oxygen, germanium and silicon shows that as long as elemental silicon or GeSi alloy is still present in the system, only silicon dioxide and elemental germanium can coexist. Any germanium dioxide present in the system will be reduced to elemental germanium by the formation of additional silicon dioxide. This process can occur only as long as silicon and germanium are mobile enough in the alloy to move with respect to each other. That is evidently the case at elevated temperatures and explains why only silicon dioxide forms. The diffusivities of germanium and silicon in GeSi have an activation energy of 3 to 5 eV [5] which is several eVs above the activation energies of the oxidation process. With failing temperature,
M.A. Nicolet, W.S. Liu /Microelectronic
Engineering 28 (1995) 185-191
187
Si Ge,Si,
_x
Figure 1. Si-Ge-0 phase diagram valid up to 1000°C (from [4]). the diffusivities of germanium and silicon in the alloy thus decrease much more rapidly than those of the oxidant in the oxide or the reaction at the interface. Ai 6OO”C, the alloy is therefore kinetically frozen, but the oxidation process is not, SC, that both silicon and germanium are oxidized as the oxidant reaches the alloy interface. This is what one observes for wet oxidation. Since germanium also enhances the reaction rate at the interface, the growth quickly becomes transportlimited and quadratic in the duration of oxidation. The structure of the mixed GeSi dioxide is amorphous and contains bonds characteristic of SiO2 in tetrahedral coordination, and of Ge02 in both tetrahedral (four-fold) and octahedral (six-fold) coordination [6]. The dry oxidation of GeSi in the low temperature regime has not been studied yet because the reaction is quite slow, but it is clear that at sufficiently low temperatures germanium will be oxidized as well. The selective oxidation of elements in an alloy occurs whenever a strong non symmetry exists in their free energy change upon oxidation, when the nonoxidizing elements can diffuse easily into the bulk, and when their segregation coefficient is large. Examples are fairly common, such as transition metal silicides, CuNi alloys, or amorphous metallic ZrAINi. At low temperatures, when the growth is parabolic, the oxide grown on polycrystalline and pseudomorphic GeSi have the same thickness (same rate constant B). The microstructure of, and the strain in, the GeSi alloy thus leave the oxidation process unchanged. This fact is consistent with the finding that the oxidation process is controlled by the transport across the oxide. Reducing the content of germanium in the alloy shifts the onset of germanium rejection to lower temperature. Going from an epitaxial to a polycrystalline film does the same thing because grain boundaries enhance diffusivities. The instability of GeSi02 in the presence of pure Si or of GeSi alloy that the phase diagram of Fig. 1. predicts is experimentally verifiable even at room temperature. After several months, crystalline precipitates appear at the oxide/alloy
188
M.A. Nicolet, W.S. Liu / Microelectronic Engineering 28 (1995) 185191
Figure 2. Cross-sectional transmission electron that micrographs show that a GeSi02 oxide is unstable in non with contact oxidized GeSi alloy (from [a]).
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interface (Fig. 2.). The precipitates have a lattice constant that is similar to that of the underlying GeSi alloy. This reaction is quite rapid at elevated temperatures (Fig. 2.). An instability of this kind is unacceptable from a device point of view and means that the gate of GeSi MOSFETs may not be formed at low temperatures simply by wet oxidation of the alloy. Dry oxidation at high temperatures does produce a stable alloyBi02 interface, but the density of interfacial states is too high for practical MOSFETs. Seen from this vantage point, the oxidation of GeSi alloy holds no promise for MOS device applications in the way the direct oxidation of pure Si does. The unstable nature of GeSi02 offers interesting opportunities for controlled phase transformation, however. Fig. 3 shows a schematic of a process sequence where an epitaxial layer of a GeSi alloy on (lOO)Si is first oxidized in a wet ambient at 700°C. This first step of oxidation is terminated before the alloy is fully consumed by the oxidation reaction. The second step is a reduction in forming gas (95% N2 + 5% H2). What happens during this reduction step depends on the concentration of Ge. If the alloy and the oxide are Ge-rich, the oxide transforms into a trilayered configuration of Si02 on top, epitaxial Ge below, and a mixed Ge and Si02 layer at the bottom, adjacent to the alloy substrate (see right-most sketch of Fig. 3.). This latter layer (the “percolation” layer) is notable in two respects: (i) it is formed largely on account of the alloy substrate during the reduction reaction, and (ii) it consists of an interconnected network of single-crystalline Ge walls and pillars that reach (“percolate”) around Si02 inclusions and establish continuous paths between the underlying alloy and the overlying Ge layer. This percolation layer is itself subdivided into distinct sublayers, as the cross-sectional micrograph of Fig. 4. shows. The figure also displays one of two remarkable properties of the Ge layer: it has fewer dislocations than the GeSi alloy seed on which it grows (about 109 versus 10 ‘0 dislocations/cm2). In addition, the Ge layer also is elastically fully relaxed, i.e. fully decoupled from its alloy substrate [7].
Figure 3. Schematic of the oxidation-reduction process sequence. The changes in the layers are approximately to scale (from [7]).
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Engineering 28 (1995) 185-191
Figure 4. Cross-sectional transmission electron micrograph of an epitaxial Geo_+i0.18 alloy after having been subjected to the process sequence of Fig. 3 (from [7]). The top SiO2 layer has been chemically removed prior to the crosssectioning of the sample. When the alloy and the GeSi02 are not rich in Ge, a percolation layer does not form, presumably because not enough germanium is available from the reduction process to form a connected network of Ge in SiO2. Rather, the reduction process results in the nucleation of nanocrystalline Ge precipitates distributed uniformly throughout the oxide layer that now consists of SiO2. The nanocrystillites are photoluminescent. The percolation layer does not form either, even with a high concentrations of Ge in the alloy and in the oxide, if pure hydrogen is used in the reduction step instead of forming gas. The strong driving force provided by pure hydrogen ambient results again in a homogeneous nucleation of Ge precipitates. It is the heterogeneous nucleation of Ge at the oxide/alloy interface from a germanium-rich alloy and oxide that favor the development of a percolation layer and thereby the growth of a relaxed epitaxial Ge layer whose crystalline quality is superior to that of the alloy from which it originates. As the structure of the percolation layer reveals, the details of the reduction mechanisms must be complex. This structure clearly results from simultaneous and competing reactions that occur in the early stages of the reduction process. The reduction of GeO2 to Ge by hydrogen near the oxide/alloy interface creates vapor which can oxidize additional alloy that itself can be reduced again by hydrogen, and so on. In effect, hydrogen acts as a catalyst that transfers oxygen from Ge02 to Si
M.A. Nicolet, W.S. Liu / Microelectronic Engineering 28 (199.5) 185191
191
deeper into the alloy and thereby creates SiO2 and Ge side by side. The speed at which these competing reactions progress must affect the morphology of the percolation layer. In parallel to this interfacial process, there must be processes that take place in the bulk of the GeSi dioxide that account for the massive relative displacements of Ge and SiO2. A dominant reaction there is most likely the partial reduction of Ge02 to GeO, which is volatile. Evidence for its presence is also found in the loss of Ge to the ambient that is observed at the beginning of the reduction process. As germanium is lost to the ambient, a dense layer of pure SiO2 begins to form near the surface. This change at the surface can explain why the loss of germanium ceases after a while. The transport of germanium as GeO in the GeSi02 also promotes the precipitation of germanium at the inter-face with the percolation layer. The growth is seeded by the germanium in the percolation layer. As the epitaxial germanium layer expands into a continuous film, the transport of gases accross the percolation layer slows down, and so does the growth of that layer. Evidently, the percolation layer is able to block the propagation of misfit dislocations across it by virtue of the discontinuous second phase of SiO2 islands that is distributed in the germanium matrix, since most of the misfit dislocations in the alloy terminate there. The scenario depicted here is largely conjectural. It will take additional investigations to substantiate or refute it. The potential payoff seems worth it, though, because the ability to improve the crystalline quality of an epitaxial layer over that of its seed is a remarkable phenomenon. Also, the growth of a highquality unstrained single-crystalline film of germanium on silicon is a capability that would permit the combination of integrated logic circuits in silicon and integrated optical systems in GaAs on the same silicon wafer. Research on GeSi is supported by the Semiconductor Research Corporation under a coordinated program between Caltech (95SJ-100) and UCLA (95SJ088), a support that is here gratefully acknowldged. We also express our gratitude to Integrated Device Technology, Inc.,San Jose, CA, for its generous grants.
References 1. B.E. Deal, A.S. Grove, J. Appl. Phys. 36 (1965) 770. 2. F.K. LeGoues, R. Rosenberg, T. Nguyen, F. Himpsel, B.S. Meyerson, J. Appl. Phys. 65 (1989) 1724. 3. R.R. Razouk, L.N. Lie, B.E. Deal, J. Electrochem. Sot. 128 (1981)2214. 4. D.C. Paine, C. Caragianis, A.F. Schwartzman, J. Appl. Phys. 70 (1991) 5076. 5. W.S. Liu, G. Bai, M.-A. Nicolet, C.H. Chern, V. Arbet, K.L. Wang, Mat. Res. Sot. Symp. Proc. Vol. 220 (1991) 259. 6. W.S. Liu, E.W. Lee, M-A. Nicolet, V. Arbet-Engels, K.L. Wang, N.M. Abuhadba, C.R. Aita, J. Appl. Phys. 71 (1992) 4015. 7. W.S. Liu, J.S. Chen, D.Y. Lie, M.-A. Nicolet, Appl. Phys. Lett. 63 (1993) 1405. 8. W.S. Liu, J.S. Chen, M.-A. Nicolet, V. Arbet-Engels, K.L. Wang, J. Appl. Phys. 72 (1992) 4444.